Correlation between microstructure and martensitic transformation, mechanical properties and elastocaloric effect in Ni–Mn-based alloys

Intermetallics 113 (2019) 106579

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Correlation between microstructure and martensitic transformation, mechanical properties and elastocaloric effect in Ni–Mn-based alloys

T

Xiao-Ming Huanga, Lu-Da Wanga, Hao-Xuan Liua, Hai-Le Yana,∗, Nan Jiaa, Bo Yanga, Zong-Bin Lia, Yu-Dong Zhangb,c, Claude Eslingb,c, Xiang Zhaoa,∗∗, Liang Zuoa a Key Laboratory for Anisotropy and Texture of Materials (Ministry of Education), School of Material Science and Engineering, Northeastern University, Shenyang 110819, China b Université de Lorraine, CNRS, Arts et Métiers ParisTech, LEM3, F-57000 Metz, France c Laboratory of Excellence on Design of Alloy Metals for low-mAss Structures (DAMAS), Université de Lorraine, 57045 Metz, France

A R T I C LE I N FO

A B S T R A C T

Keywords: Magnetic shape memory alloys Ni–Mn-Based alloys Elastocaloric effect Martensitic transformation Microstructure Directional solidification.

Microstructure is a crucial parameter to affect elastocaloric effect in the Ni–Mn-based alloys. However, the specific influence of different microstructural parameters on elastocaloric effect remains unclear. In the present work, the effect and underlying mechanism of microstructure, including grain size and texture, on the martensitic transformation, mechanical properties and elastocaloric effect in Ni50Mn34.8In15.2 alloy were studied. Three fabrication techniques including the arc-melting, suction casting and directional solidification were utilized to obtain different microstructures. Results show that the phase transition temperature and transformation entropy change are insensitive to the microstructure. While, the microstructure has an obvious influence on phase transition interval ΔTInt. The suction casting sample has the widest ΔTInt. Moreover, we observed that the adiabatic temperature change ΔTad largely depended on the feature of microstructure. The grain boundary is disadvantageous to ΔTad. On the one hand, it would interact with habit plane and bring about energy consumption. On the other hand, due to the enhanced energy for full transformation, it would reduce the volume fraction of phase transition under certain external strain or force. The fracture mechanism of the Ni–Mn-based alloy is independent of microstructure, i.e. brittle fracture with intergranular pattern. It is attributed to the strong p-d covalent hybridization between Ni and main-group element (In). Strong texture and coarse grain are favorable to improve the mechanical properties. This study may provide some guidance for microstructure design to improve the elastocaloric effect for the Ni–Mn-based alloys.

1. Introduction Due to the environmental friendliness, high energy efficiency and promising potential to replace traditional vapor compression technology, the solid-state refrigerant based on elastocaloric effect has become a hot topic in both material science and condensed matter physics [1–8]. Among these materials, Ni–Mn-based alloys are attracting more and more attention due to their low critical stress and large specific adiabatic temperature change (ΔTad/σ or ΔTad/ε) [9–19], which is greatly desirable for the applications in the micro- or nano-refrigeration fields, such as microelectronic chip. Apart from mechanical filed, the giant caloric effect can be realized by applying magnetic field, i.e. magnetocaloric effect. On the one hand, the combination of these effects can greatly increase the refrigeration efficiency. On the other

∗

hand, as evidenced in Refs. [20,21], the multi-excitation of caloric effect in Ni–Mn–In alloys supplies us an opportunity to reduce the hysteresis loss and further improve the cyclic stability. Unfortunately, the relatively poor ductility and low cycle stability of these materials greatly hinder their practical application [22–24]. To improve the mechanical properties of Ni–Mn-based alloys, a general strategy is alloying to tune the electronic structure or introduce the soft second phase by substituting/adding the fourth (even fifth) elements, such as Cu [25,26], Cr [27], Ti [28,29] and Fe [30,31]. Different from numerous physical properties that are majorly decided by alloy composition, the elastocaloric effect, strongly depended on the mechanical properties of materials, is closely related to the microstructure. This provides us an extra dimension, i.e. microstructure design, to improve it. However, the effect and underlying mechanism of microstructure on

Corresponding author. Corresponding author. E-mail addresses: [email protected] (H.-L. Yan), [email protected] (X. Zhao).

∗∗

https://doi.org/10.1016/j.intermet.2019.106579 Received 15 April 2019; Received in revised form 9 July 2019; Accepted 8 August 2019 Available online 20 August 2019 0966-9795/ © 2019 Elsevier Ltd. All rights reserved.

Intermetallics 113 (2019) 106579

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homogenized at 1173 K for 24 h followed by quenching into water. The behaviors of martensitic transformation were examined by differential scanning calorimetry (DSC, TA-Q100). The morphology and composition of the samples were investigated by scanning electron microscopy (SEM, Jeol JSM 6500F) with an X-ray energy dispersive spectroscopy (EDS) analysis system. Three incomplete pole figures of {220}A, {224}A and {004}A were measured by the Schulz back-reflection method using Cu-Kα radiation in a Rigaku SmartLab X-ray diffractometer. The complete pole figures of {220}A, {224}A and {004}A were then calculated by using Mtex software package [35]. The compressive experiment was conducted on a universal testing machine (AGXPlus 100 kN) at room temperature. Rectangular samples with the size of 5 mm × 3 mm × 8 mm were used. The temperature variation during the stress-induced martensitic transformation was recorded by a thermocouple attached to the sample surface. The Vienna ab initio simulation package (VASP) was used to calculate the electronic localization function (ELF) of Ni8Mn6In2 alloy. A kinetic energy cutoff of 600 eV, energy convergence criterion of 10–5 eV and 15 × 15 × 15 k-points were adopted. The valence configurations of Ni: 3d84s2, Mn: 3d54s2 and In: 5s25p1 were taken, respectively.

mechanical properties and elastocaloric effect for the Ni–Mn-based alloys remain unclear. Numerous studies have shown that the grain size and crystallographic preferred orientation (texture) are one of the most important microstructure parameters to affect the mechanical and elastocaloric behaviors in the Ni–Mn-based alloys [29,32,33]. For the effect of grain size, V. Sánchezalarcos et al. [29] observed that the microhardness and yield strength of (Ni50Mn33.5In16.5)100-xTix (x = 0, 0.5, 1, 2) alloys can be largely enhanced with the fine grain caused by Ti doping; Z. Yang et al. [32] found that the grain refinement can improve the fatigue life of elastocaloric effect in the suction casting (Ni51.5Mn33In15.5)99.7B0.3 alloy. On the contrary, recently, Z. Li et al. [33] suggest that the coarse grain can improve the mechanical properties in Ni45.7Co4.2Mn37.3Sb12.8 alloy. As for the effect of texture, Y. J. Huang et al. [34] reported that the textured microstructure can promote the mechanical properties of Ni48Mn35In17 alloy. A prominent adiabatic temperature change ΔTad of about −4 K was obtained in this alloy. D. W. Zhao et al. [26] suggested that the texture is favorable to elevate the ΔTad by comparing the highly textured and non-textured sample in Ni50Mn31.5In16Cu2.5 alloy. As summarised above, it is clear now that the microstructure is a crucial parameter to affect the mechanical properties and elastocaloric effect for the Ni–Mn-based alloys. However, the specific effects of microstructure on the mechanical properties and elastocaloric effect, as well as the behaviors of martensitic transformation, are still unclear. For example, the effect and underlying mechanism of microstructure on ΔTad of elastocaloric effect remain unknown. To our knowledge, under the same measurement condition (strain and strain rate), the ΔTad for the samples with the same composition but different microstructure have never been reported. Moreover, in the literature, both positive and negative effect of grain size on mechanical and elastocaloric behaviors have been suggested [29,32,33]. In our view, one of the reasons might be the influence of alloying element. When the microstructure variation is induced by alloying element, such as the grain refinement induced by Ti doping [29], it is difficult to distinguish the specific influence of these two factors. As a result, a specialized study on the influence of microstructure on mechanical and elastocaloric effect is in great need. The purpose of this work is to reveal the effect and underlying mechanism of microstructure, specifically grain size and texture, on mechanical properties and elastocaloric effect, as well as on the behaviors of martensitic transformation, for the Ni–Mn-based alloys. Three kinds of extensively used fabrication methods including the arcmelting, suction casting and directional solidification were exploited so as to obtain the samples with different microstructure. To exclude the influence of alloying element, the same alloy with nominal composition of Ni50Mn34.8In15.2 (at. %) was chosen for different preparation methods. Ab-initio calculation method was used to study the features of chemical bonding to reveal the nature of intrinsic brittleness of these compounds. This work could be expected to provide some instructions for microstructure design to optimize the mechanical properties and elastocaloric effect for the Ni–Mn-based alloys.

3. Results and discussion 3.1. Microstructure and martensitic transformation Figs. 1a1, b1 and c1 display the backscattered electron (BSE) images for the Ni50Mn34.8In15.2 samples prepared by means of the arc-melting, suction-casting and directional solidification techniques, respectively. For all samples, we find that the alloy is in a single austenite phase at room temperature. As seen in Fig. 1a1, the arc-melting sample possesses columnar-shape grains with the long axis of 1–2 mm and the wide axis of 100–500 μm. By means of the suction casting technique, the equiaxed grains with grain size less than 200 μm is obtained owing to the ultrahigh cooling rate of this method. After the process of directional solidification, the sample has coarse columnar-shaped grains (several mm) along the direction of directional solidification (DS). In Figs. 1a2, b2 and c2, we display {004}A complete pole figures for the arc-melting, suction casting and directional solidification samples, respectively. It is seen that the samples prepared by arc-melting and suction-casting techniques show very weak textures, i.e., no remarkable crystallographic preferred orientation. On the contrary, as seen in Fig. 1c2, the directional solidification sample exhibits a strong cube texture with the [001]A//Z0 (the direction of DS). It is because the grains of directional solidification sample generally prefer to grow with their easiest growth direction, i.e. [001]A for cubic crystals, along the direction with the largest temperature gradient (Z0). The obtained bulk-type texture (rather than fiber-type texture) could be due to the huge grain size of sample prepared by directional solidification and additional heat treatment. The different microstructures of these three samples allow us to explore the influence of microstructure feature on martensitic transformation, mechanical properties and elastocaloric effect for the Ni–Mn-based alloys. Fig. 2 shows the DSC curves for different Ni50Mn34.8In15.2 samples. With the tangent method, the critical start and finish temperatures of the forward and reverse martensitic transformation, i.e., Ms, Mf, As and Af, were determined, as listed in Table 1. For different samples, we find that their Ms are almost the same in consideration of measurement error. As a result, the chemical composition discrepancy caused by different fabrication methods could be negligible, since the phase transition temperature of Ni–Mn-based alloys is very sensitive to the alloy composition [20]. This is further confirmed by the EDS measurements, as seen in Table 1. On the contrary, the phase transition intervals ΔTInt (Ms–Mf) for different samples have an obvious discrepancy. The suction cast sample, as seen in Fig. 2, has the widest phase transition interval (~6.8 K). This could be ascribed to the largest fraction of crystal defects (grain boundaries) of this sample brought

2. Experimental details The alloy ingots with nominal composition of Ni50Mn34.8In15.2 (at. %) were prepared by the arc-melting technique in a high-purity argon atmosphere using the high-purity raw elements (the arc melting sample). Considering the high volatility of Mn in comparison to other constituents, an excess 1 wt % of Mn was added to compensate the loss during melting. Part of the ingots were remelted and ejected into a copper mold with the diameter of 10 mm under the protection of argon (the suction casting sample). To obtain the sample with a coarse grain and a strong texture, part of the rods was directionally solidified by using the liquid-metal-cooling technique (the directional solidification sample). The growing rate is set to be 0.05 mm/s. To exclude the influence of atom disorder, all the arc melting, suction casting and directional solidification samples were sealed into quartz tubes and 2

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Fig. 1. Backscattered electron images and {004}A complete pole figures for the Ni50Mn34.8In15.2 samples prepared by (a1 and a2) the arc-melting, (b1 and b2) suction casting and (c1 and c2) directional solidification (DS) techniques. The solid arrows in (a1, b1 and c1) indicate the compression directions (CD) for the following mechanical and elastocaloric property measurements. The direction of DS in (c1) is indicated by the open arrow. In (c2), Z0 is the direction of solidification DS; X0 is an arbitrary direction perpendicular to the direction of DS; Y0 is the cross product of Z0 by X0.

evaluated by ΔH/T0 in which ΔH represents the enthalpy change determined by integrating the area of exothermic peak of DSC curve and T0 is equilibrium temperature of martensitic transformation defined as (Ms + Mf)/2, as listed in Table 1. It is seen that the arc-melting, suction casting and directional solidification samples have similar ΔSTr, i.e., 16.7, 16.4 and 17.7 J kg−1 K−1, respectively. It shows that the ΔSTr of martensitic transformation is insensitive to the grain size and crystallographic preferred orientation. The slightly smaller ΔSTr for the suction casting sample might be related to the more energy consumption caused by the interactions between habit plane and grain boundaries during martensitic transformation. As for the thermal hysteresis ΔTHyst ((As + Af)/2–(Ms + Mf)/2), we find that all these three samples have also similar values (around 13 K) with a difference less than 1 K. As is well known, the ΔTHyst is intrinsically originated from the geometric incompatibility between austenite and martensite [36,37], which can be quantitatively represented by the middle eigenvalue (λ2) of the stretch tensor and the norm of XI for the type-I twin and XII for the typeII twin. In this work, the geometric incompatibilities between austenite and martensite for these three samples should be similar due to their negligible composition difference (Table 1), which accounts for their similar ΔTHyst. Fig. 2. DSC curves for different Ni50Mn34.8In15.2 samples. Ms, Mf, As and Af indicate the start and finish temperatures for the forward and reverse martensitic transformation, respectively. The testing temperature (TTest) for measuring elastocaloric effect is indicated by the dashed arrow. The endothermic peak (Endo) represents the reverse martensitic transformation.

3.2. Mechanical properties Fig. 3a displays the compressive fracture curves for different Ni50Mn34.8In15.2 samples with a strain rate of 1 × 10−4 s−1. The compression directions (CD) for various samples are indicated in Figs. 1a, b and c, respectively. For all samples, it is seen that the compressive fracture curve can be divided into four stages: AB, BC, CD and DE. The AB stage, a linear relation between strain and stress, represents the elastic deformation of austenite. When the load reaches to the critical stress (σCrit) of martensitic transformation, an obvious strain-

about by relatively small grain size, which might act as roadblocks to impede the movement of habit plane and thus delaying the process of martensitic transformation. By contrast, the directional solidification sample has the narrowest ΔTInt (~4.6 K). The transformation entropy change ΔSTr for different samples were

Table 1 Martensitic transformation start (Ms) and finish (Mf) temperatures, reverse martensitic transformation start (As) and finish (Af) temperatures, phase transition interval ΔTInt (Ms–Mf), thermal hysteresis ΔTHyst ((As + Af)/2–(Ms + Mf)/2), transformation entropy change (ΔSTr) for different Ni50Mn34.8In15.2 samples. Fabrication methods

Actual composition

Ms (K)

Mf (K)

As (K)

Af (K)

ΔTInt (K)

ΔTHyst (K)

ΔSTr (J kg−1 K−1)

Arc-melting Suction casting Directional solidification

Ni50.0Mn34.5In15.5 Ni49.6Mn34.8In15.6 Ni49.6Mn34.7In15.7

266.4 265.6 267.8

260.7 258.8 263.2

272.6 272.3 274.6

280.2 279.4 281.8

5.7 6.8 4.6

12.9 13.7 12.7

16.7 16.4 17.7

3

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Fig. 3. (a) Compressive fracture curves for different Ni50Mn34.8In15.2 samples measured at room temperature. The σComp, εComp, σCrit and εTrans indicate the compressive strength, compressive strain, critical stress to induce martensitic transformation and transformation strain during martensitic transformation, respectively. Fracture surfaces for the samples fabricated by (b) the arc-melting, (c) the suction casting and (d) the directional solidification technique. (e) Distribution of electron localization function (ELF) on (110) plane of austenite.

discussed later.

stress plateau originated from the stress-induced martensitic transformation, i.e., the BC stage, is observed. In the CD stage, the sample mainly undergoes the elastic deformation of martensite. When the load is larger than compressive strength (σComp), i.e., the DE stage, the crack nucleates and propagates with the increasing stress, which finally results in the fracture of the sample. Comparing the compressive fracture curves for different samples, we find that the arc-melting sample has a medium compressive strength σComp (~451 MPa) and compressive strain εComp (~9.4%). After the suction casting, unexpectedly, the σComp and εComp decrease to ~341 MPa and ~7.0%, respectively, although the grain size is smaller than that prepared by arc-melting method (Fig. 1). For the directional solidification sample, both σComp and εComp are remarkably enhanced to ~991 MPa and ~11.9%, respectively. It clearly shows that the directional solidification technique can greatly improve the fracture resistance. Moreover, the behaviors of stress-induced martensitic transformation for different samples have also large discrepancy. With the tangent method, the σCrit and transformation strain (εTrans) for the arcmelting sample are measured to be 112 MPa and 3.3%, respectively. After the suction casting, the σCrit slightly increases to 149 MPa. This could be due to the strongly constrained effect between neighboring grains with different orientations during the stress-induced martensitic transformation [26]. On the contrary, the σCrit of the directional solidification sample reduces to 72.6 MPa, which is greatly desirable in view of machine design. Meanwhile, the εTrans of this sample is enhanced to 5.7%. To uncover the underlying mechanism of microstructure effect on the mechanical properties, the morphologies of fracture surface for different samples were investigated, as shown in Figs. 3b, c and d. For all the samples, we observe that the fracture surfaces exhibit typical brittle characteristics with intergranular pattern. This is greatly different from the conventional structurally metallic materials, such as steel, in which the fracture behavior is significantly depended on the microstructure [38,39]. In order to reveal the nature of the microstructure-insensitive fracture behavior of the studied alloy, the electron localization function (ELF) was calculated by means of density functional theory. For simplicity, a superstructure model with the composition of Ni8Mn6In2, similarity to the studied Ni50Mn34.8In15.2 alloy, is exploited. In Fig. 3e, we plot the distribution of ELF on (110) plane of austenite. As indicated in the dashed lines, strong covalent bonding between Ni and the main-group element (In), related to the p-d orbital hybridization [40,41], are observed. Different from the metallic bond, the covalent bond possesses high binding energy and is strongly directional, e.g., the < 111 > direction in the Ni–Mn-based alloys (Fig. 3e). Hence, the nucleation and propagation for dislocations should be difficult since their related bond distortion and breaking generally requires huge energy. This might be the reason for the microstructureinsensitive fracture behavior in the Ni–Mn-based alloys. The mechanism for the effect of microstructure on σComp and εComp will be

3.3. Elastocaloric effects We now turn our attention to the influence of microstructure on elastocaloric effect. During measurements, the samples were first compressed to 7% with a strain rate of 2.1 × 10−3 s−1; then, held on 40 s to return its initial temperature; finally, unloaded at a strain rate of 3.1 × 10−2 s−1 for adiabatic unloading process. This loading strategy has been extensively utilized to study the elastocaloric refrigerants, such as Ni–Mn-based alloys [23,34,42], Co-based [43], Cu-based [44] and Ni–Ti [45] alloys. Fig. 4 shows the temperature variations across the stress-induced (reverse) martensitic transformation for the samples prepared by different methods. It is seen that the adiabatic temperature change ΔTad for the arc-melting, suction casting and directional solidification samples are −4.6 K, −3.6 K and −7.6 K, respectively. It clearly shows that the microstructure of material plays a critical role on ΔTad of elastocaloric effect. In Table 2, we compare the ΔTad value of the directional solidification sample with some other Heusler-type alloys reported in the literature. It is seen that the ΔTad obtained in this work is comparable to the highest elastocaloric effects of polycrystalline or even single crystalline materials in the Ni–Mn-based alloys. The specific adiabatic temperature change |ΔT/σ| for the directionally solidified sample is determined to be 22.0 K/GPa, which is larger than that of Ti54.9Ni32.5Cu12.6 (~17.1 K/GPa) [46] and Fe68.8Pd31.2 (~20.0 K/GPa) [47] alloys. As is well known, the ΔTad during the stress-induced martensitic

Fig. 4. Temperature-time profiles of Ni50Mn34.8In15.2 samples prepared by different methods. 4

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Table 2 Comparison of adiabatic temperature change ΔTad of the present directionally solidified Ni50Mn34.8In15.2 sample and some other Heusler-type alloys reported in the literature. Alloy composition (at. %)

Sample state

ΔTad (K)

Reference

Ni45Mn36.4In13.6Co5 Ni45Mn36.2In13.8Co5 Ni45Mn36In13Co5Cr (Ni51.5Mn33In15.5)99.7B0.3 Ni51.4Mn34In15.6Tb0.4 Ni48Mn35In17 Ni45.7Mn36.6In13.3Co5.1 Ni50Mn31.5In16Cu2.5 Ni57Mn18Ga21In4 Ni50Mn34.8In15.2

Arc melting polycrystalline Arc melting polycrystalline Arc melting polycrystalline Suction casting polycrystalline Suction casting polycrystalline Directionally solidified polycrystalline Directionally solidified polycrystalline Directionally solidified polycrystalline Single crystal Directionally solidified oligocrystal

−4 −3.1 −5.8 −6.6 −5.1 −4 −3.5 −10 −9.6 −7.6

[22] [48] [27] [32] [10] [34] [49] [26] [50] This work

elastocaloric effect. As illustrated above, the microstructure features including grain size and texture have important influence on the fracture resistance and cycle stability. We now discuss their possible underlying mechanisms. In view of lattice dynamics, the martensitic transformation of the Ni–Mn-based alloy is attributed to the pronounced phonon softening of the TA2- < 110 > mode, which is originated from the phenomenon of fermi surface nesting under the strong electron-phonon coupling [51,52]. The low energy of TA2- < 110 > mode results that the phase transition is majorly realized by shearing {110} plane along < −110 > direction, as shown in Fig. 6a, to reduce the energy of the system. It should be noted that the transformation system of {110} < −110 > is the only activated deformation system to accommodate external strain during the stress-induced martensitic transformation. Moreover, it is known that the crack is resulted from the strain discontinuity in local region of materials when the resolved stress is larger than the facture stress. For the Ni–Mn-based alloys, the failure under mechanical load is mainly attributed to the strain incompatibility near grain boundaries (Fig. 3). Thus, we will discuss, respectively, the influence of texture and grain size on the strain compatibility near grain boundaries under the shear-type {110} < −110 > transformation systems. In Figs. 6b and c, we show the effect of texture on the strain compatibility near grain boundary. Grains A and B (Fig. 6b) with similar orientation and Grains C and D (Fig. 6c) with different orientation represent the strong and random texture, respectively. The lattices depicted in solid and dashed lines indicate austenite and martensite, respectively. For the material with strong texture, i.e. the adjacent grains with similar orientation (Fig. 6b), the transformation strains between adjacent grains would be compatible near grain boundary. By contrast, when the adjacent grains have a large misorientation (random texture, Fig. 6c), the strain discontinuity would inevitably occur near grain boundary, especially at triple junctions. It would result in the nucleation and propagation of crack along the grain boundary due to the inherent weakness of grain boundary relative to grain interior. Thus, the strong texture should be favorable to increase the fracture resistance and cyclic stability for the Ni–Mn-based alloys. We then discuss the influence of grain size. As demonstrated above, the fracture of the Ni–Mn-based alloys is mainly ascribed to the strain incompatibility near grain boundaries. The fine grain generally results in a higher density of grain boundaries. Therefore, the nucleation sites for crack in materials with fine grains should be larger than that in materials with coarse grains. As seen in Fig. 1, the suction casting sample has the smallest grain size and random texture. Both of them are unfavorable to the fracture resistance, which accounts well for its worst mechanical properties. On the contrary, both these microstructural parameters in the directional solidification sample, i.e., coarse grain and strong texture, are favorable to the mechanical properties. This explains its good fracture resistance and cyclic stability. Moreover, it should be noted that the cycle stability for the state-of-the-art directionally solidified Ni–Mn-based alloys [11,53] still cannot meet the practical applications (~107). Apart from the microstructure

transformation can be estimated by the following relation [48–50]: ΔTM→A = ΔTvib + ΔTmag +ΔTelec + ΔTfri

(1)

where the ΔTvib, ΔTmag and ΔTelec indicate the temperature variations caused by the lattice vibration entropy change, magnetic entropy change and electronic entropy change, respectively. The ΔTfri represents the friction heat generated from the movements of habit plane as well as martensite variant interface during martensitic transformation. In this work, the sum of ΔTvib, ΔTmag and ΔTelec for different samples should be the same due to their similar chemical composition. However, owing to the high loading rate, the ΔTfri for samples with different microstructure should have large discrepancy. For the suction casting sample, owing to its substantial grain boundaries, the ΔTfri caused by the interactions between the habit plane/martensite variant interface and crystal defects might be much larger than the other samples. This could be one of reasons for its worst cooling effect. Moreover, as seen in Fig. 3a, the suction casting sample has the largest σCrit and the smallest εTrans. Therefore, in view of dynamics of phase transition, the maximum energy required to realize the complete phase transition for the suction casting sample should be larger than the others. Therefore, when applying the same force or the same strain, the volume fraction of phase transition for the suction casting sample should be smaller than the other samples. This might be another reason for the worst cooling effect of suction casting sample. Apart from the ΔTad, the cyclic stability is another paramount factor for elastocaloric effect in view of practical application. In Fig. 5, we plot the cyclic stress-strain curves and the corresponding temperature variations for various Ni50Mn34.8In15.2 samples prepared by different methods. During these measurements, the 40 loading-unloading cycles with the maximum strain of 6% at a strain rate of 8.3 × 10−3 s−1 were performed. Comparing the stress-strain curves for different samples (Figs. 5a1, b1 and c1), we find that the cyclic stability of directional solidification sample is obviously superior to the others. It is accord with the fracture behaviours of these samples (Fig. 3a). At the first cycle, the suction casting sample has a large irreversible residual strain εResi (~1.2%). On the contrary, the εResi for the arc-melting and directional solidification samples are determined to be ~0.3% and ~0.1%, respectively, which are one magnitude order smaller than that in the suction casting sample. As seen in Fig. 2, the finish temperature (Af ≈ 280 K) of martensitic transformation for the studied alloy is far below the testing temperature (TTEST ≈ 298 K). As a result, the εResi should be majorly attributed to the crystal defects (grain boundaries) rather than the irreversibly remnant martensite. The largest amount grain boundaries of the suction casting sample accounts well for its largest residual strain. Comparing the cyclic temperature variation curves, for the arc-melting and suction casting samples (Figs. 5a2 and b2), the ΔTad gradually decreases with the increasing cycle number. While, the ΔTad of directional solidification sample is hardly degradative after 40 loading/unloading cycles. It further confirms the prominent influence of microstructure on the cyclic stability of

5

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Fig. 5. Stress-strain and temperature-time curves during cyclic loading, holding and unloading processes for Ni50Mn34.8In15.2 samples prepared by the (a1 and a2) arcmelting, (b1 and b2) suction casting and (c1 and c2) directional solidification methods. All measurements are carried out with a maximum strain of 6% at a strain rate 8.3 × 10−3 s−1 at room temperature.

modification, other strategies to enhance the mechanical properties, such as grain boundary strengthening, should also be considered in the future study. In addition, the combination of multicaloric effects by applying different external fields (thermal, mechanical and magnetic fields) simultaneously or successively is an effective strategy to reduce the hysteresis loss and further increase the cyclic stability. 4. Conclusions In the present work, the influence of microstructure features including grain size and texture on the martensitic transformation, mechanical behavior and elastocaloric effect for the polycrystalline Ni50Mn34.8In15.2 alloys were studied. Three different fabrication methods, i.e., the arc-melting, suction casting and directional solidification, are considered. The main conclusions are as follows: (1) The martensitic transformation temperature and transformation entropy change ΔSTr are insensitive to the microstructure. While, the microstructure has an obvious influence on the phase transition interval ΔTInt. The suction casting sample has the widest ΔTInt, which might be due to its largest amount grain boundaries that could act as roadblocks to impede the movement of habit plane and then to delay the process of martensitic transformation. (2) The microstructure has an important effect on the elastocaloric effect. The crystal defects, such as grain boundaries, is disadvantageous to the adiabatic temperature change ΔTad. On the one hand, it would interact with the habit plane/martensite variant interface and bring about the energy consumption. On the other hand, it would reduce the volume fraction of phase transition under certain external strain or force. (3) The facture surfaces of all samples show typical brittle characteristics with intergranular patterns owing to the strong p-d covalent

Fig. 6. (a) {110} < −110 > transformation system for the Ni–Mn-based alloys. Illustration of the effects of (b) strong texture and (c) random texture on the strain compatibility near grain boundary.

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hybridization between 3d electrons of Ni and p electrons of In. The strong texture is favorable to the fracture resistance and cyclic stability, since it can reduce the strain incompatibility near grain boundaries. The suction casting sample possesses the poorest mechanical properties owing to its largest amount grain boundaries that are beneficial to the crack nucleation.

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Acknowledgements This work is supported by the National Natural Science Foundation of China (Grant No. 51431005, 51801020, 51771044), the Fundamental Research Funds for the Central Universities (Grant No. N170203009), the Liao Ning Revitalization Talents Program (Grant No. XLYC1802023), the PhD Starting Foundation of Liaoning Province (Grant No. 20180540115) and Postdoctoral Science Foundation of China (Grant No. 2018M641700). References [1] E. Bonnot, R. Romero, L. Mañosa, E. Vives, A. Planes, Elastocaloric effect associated with the martensitic transition in shape-memory alloys, Phys. Rev. Lett. 100 (12) (2008) 125901. [2] J.L. Sánchez Llamazares, H. Flores-Zuniga, C. Sanchez-Valdes, C.A. Ross, C. Garcia, Refrigerant capacity of austenite in as-quenched and annealed Ni51.1Mn31.2In17.7 melt spun ribbons, J. Appl. Phys. 111 (7) (2012) 07A932. [3] L. Manosa, S. Jarque-Farnos, E. Vives, A. Planes, Large temperature span and giant refrigerant capacity in elastocaloric Cu-Zn-Al shape memory alloys, Appl. Phys. Lett. 103 (21) (2013) 211904. [4] H. Ossmer, F. Lambrecht, M. Gueltig, C. Chluba, E. Quandt, M. Kohl, Evolution of temperature profiles in TiNi films for elastocaloric cooling, Acta Mater. 81 (2014) 9–20. [5] M. Schmidt, A. Schütze, S. Seelecke, Scientific test setup for investigation of shape memory alloy based elastocaloric cooling processes, Int. J. Refrig. 54 (2015) 88–97. [6] F. Chen, Y.X. Tong, L. Li, J.L. Sánchez Llamazares, C.F. Sanchez-Valdes, P. Mullner, Broad first-order magnetic entropy change curve in directionally solidified polycrystalline Ni-Co-Mn-In, J. Alloy. Comp. 727 (2017) 603–609. [7] Z.B. Li, Y.W. Jiang, Z.Z. Li, C.F. Sanchez Valdes, J.L. Sanchez Llamazares, B. Yang, Y.D. Zhang, C. Esling, X. Zhao, L. Zuo, Phase transition and magnetocaloric properties of Mn50Ni42-xCoxSn8 (0≤x≤10) melt-spun ribbons, IUCr J. 5 (1) (2018) 54–66. [8] H.L. Yan, C.F. Sánchez-Valdés, Y. Zhang, J.L. Sánchez Llamazares, Z. Li, B. Yang, C. Esling, X. Zhao, L. Zuo, Correlation between crystallographic and microstructural features and low hysteresis behavior in Ni50.0Mn35.25In14.75 melt-spun ribbons, J. Alloy. Comp. 767 (2018) 544–551. [9] P.O. Castillo-Villa, L. Manosa, A. Planes, D.E. Soto-Parra, J.L. Sanchez-Llamazares, H. Flores-Zuniga, C. Frontera, Elastocaloric and magnetocaloric effects in Ni-Mn-Sn (Cu) shape-memory alloy, J. Appl. Phys. 113 (5) (2013). [10] Q. Shen, D.W. Zhao, W. Sun, Y. Li, J. Liu, The effect of Tb on elastocaloric and mechanical properties of Ni-Mn-In-Tb alloys, J. Alloy. Comp. 696 (2017) 538–542. [11] Y. Hu, Z.B. Li, B. Yang, S.X. Qian, W.M. Gan, Y.Y. Gong, Y. Li, D.W. Zhao, J. Liu, X. Zhao, L. Zuo, D.H. Wang, Y.W. Du, Combined caloric effects in a multiferroic NiMn-Ga alloy with broad refrigeration temperature region, Apl. Mater. 5 (4) (2017) 046103. [12] W. Sun, J. Liu, B. Lu, Y. Li, A. Yan, Large elastocaloric effect at small transformation strain in Ni45Mn44Sn11 metamagnetic shape memory alloys, Scr. Mater. 114 (2016) 1–4. [13] T. Gottschall, A. Gracia-Condal, M. Fries, A. Taubel, L. Pfeuffer, L. Manosa, A. Planes, K.P. Skokov, O. Gutfleisch, A multicaloric cooling cycle that exploits thermal hysteresis, Nat. Mater. 17 (10) (2018) 929–934. [14] H.L. Yan, Y.D. Zhang, N. Xu, A. Senyshyn, H.-G. Brokmeier, C. Esling, X. Zhao, L. Zuo, Crystal structure determination of incommensurate modulated martensite in Ni−Mn−In Heusler alloys, Acta Mater. 88 (2015) 375–388. [15] H.L. Yan, B. Yang, Y.D. Zhang, Z.B. Li, C. Esling, X. Zhao, L. Zuo, Variant organization and mechanical detwinning of modulated martensite in Ni−Mn−In metamagnetic shape-memory alloys, Acta Mater. 111 (2016) 75–84. [16] Y.H. Qu, D.Y. Cong, S.H. Li, W.Y. Gui, Z.H. Nie, M.H. Zhang, Y. Ren, Y.D. Wang, Simultaneously achieved large reversible elastocaloric and magnetocaloric effects and their coupling in a magnetic shape memory alloy, Acta Mater. 151 (2018) 41–55. [17] D.Y. Cong, L. Huang, V. Hardy, D. Bourgault, X.M. Sun, Z.H. Nie, M.G. Wang, Y. Ren, P. Entel, Y.D. Wang, Low-field-actuated giant magnetocaloric effect and excellent mechanical properties in a NiMn-based multiferroic alloy, Acta Mater. 146 (2018) 142–151. [18] R. Chulist, L. Straka, A. Sozinov, T. Tokarski, W. Skrotzki, Branched needle microstructure in Ni−Mn−Ga 10M martensite: EBSD study, Acta Mater. 128 (2017) 113–119. [19] R. Chulist, L. Straka, A. Sozinov, T. Lippmann, W. Skrotzki, Modulation reorientation in 10M Ni-Mn-Ga martensite, Scr. Mater. 68 (9) (2013) 671–674. [20] J. Liu, T. Gottschall, K.P. Skokov, J.D. Moore, O. Gutfleisch, Giant magnetocaloric effect driven by structural transitions, Nat. Mater. 11 (7) (2012) 620–626.

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